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  • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
  • C21D8/0226—Hot rolling
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  • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
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  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/004—Dispersions; Precipitations
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  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/008—Martensite

Definitions

• low-alloy steels such as Cr-Mo steel
• the materials for 35 MPa or higher pressure hydrogen accumulators are limited to aluminum alloys and SUS316 that are less susceptible to material quality deterioration by hydrogen.
• Patent Literature 4 proposes a steel for high-pressure hydrogen environment in which fine V-Mo precipitates are used as sites for trapping diffusible hydrogen in the steel, thereby suppressing embrittlement caused by diffusible hydrogen.
• Patent Literature 5 proposes a low-alloy high-strength steel with excellent resistance to embrittlement in a high-pressure hydrogen environment.
• the tensile strength is controlled within a very narrow range of 900 to 950 MPa by tempering a Cr-Mo steel at a relatively high temperature in refining of the steel.
• a high-pressure hydrogen gas environment means an environment at a total pressure of 1 MPa or more and a hydrogen partial pressure of 1 MPa or more.
• low-alloy steels used for hydrogen pressure vessels or the like undergo hydrogen embrittlement to suffer a decrease in the tensile strength of the materials or to exhibit a significantly accelerated fatigue crack propagation rate.
• Enhancing the material strength allows pressure vessels, such as high-pressure hydrogen accumulators, to be designed with a small wall thickness.
• high-strength steel materials are increasingly needed for the purposes of increasing the storage capacity and reducing the product weight.
• a steel material is strengthened to 850 MPa or more, the hydrogen embrittlement phenomenon becomes noticeable and the steel may be fractured below the maximum tensile strength tested in the air.
• the present invention has been made in consideration of the circumstances discussed above. It is therefore an object of the present invention to provide a high-strength seamless steel pipe for high-pressure hydrogen container that has excellent hydrogen embrittlement resistance, and a method for manufacturing the same.
• the present inventors Based on the fact that a desired high strength and hydrogen embrittlement resistance need to be satisfied concurrently in order to achieve the above objective high-strength seamless steel pipe for high-pressure hydrogen container (hereinafter, also written simply as the high-strength seamless steel pipe), the present inventors carried out extensive studies on various factors that would affect strength and hydrogen embrittlement resistance. As a result, the present inventors have found that the hydrogen embrittlement resistance of a high-strength steel pipe with a tensile strength TS of 850 MPa or more is significantly affected by the presence or absence of intergranular fracture.
• the present inventors have then found that fine Mo precipitates that are dispersed in the steel microstructure serve as powerful hydrogen trapping sites to reduce hydrogen accumulation at grain boundaries and to suppress the occurrence of intergranular fracture. Furthermore, the present inventors have found that the hydrogen embrittlement resistance is enhanced with increasing amount of Mo precipitates having a size of 50 nm or less.
• the present inventors have come to the realization that a high-strength seamless steel pipe will be further enhanced in hydrogen embrittlement resistance by increasing the amount of molybdenum contained in precipitates and, in particular, controlling the amount of molybdenum contained in 50 nm and smaller precipitates to or above an appropriate level.
• a high-strength seamless steel pipe for high-pressure hydrogen container that has a high tensile strength TS of 850 MPa or more and excellent hydrogen embrittlement resistance can be manufactured easily and inexpensively.
• the present invention thus achieves significant industrial effects.
• the manufacturing method involves appropriate amounts of appropriate alloying elements and promotes the formation of Mo precipitates. In this manner, the manufacturing method can stably produce high-strength seamless steel pipes that have a high strength desired for pressure vessels and also have excellent hydrogen embrittlement resistance.
• composition of the high-strength seamless steel pipe for high-pressure hydrogen container (hereinafter, also written simply as the high-strength seamless steel pipe) of the present invention is limited.
• mass% in the composition is simply written as %.
• Carbon contributes to increasing the strength of steel by solid solution strengthening and enhances the hardenability of steel so as to contribute to the formation of a microstructure mainly composed of a martensite phase at the time of quenching.
• the C content needs to be 0.20% or more.
• the C content is preferably 0.22% or more, more preferably 0.25% or more, and still more preferably 0.28% or more. If, on the other hand, the C content is more than 0.50%, cracking occurs during quenching to significantly deteriorate the manufacturability. Thus, the C content is limited to 0.50% or less.
• the C content is preferably 0.45% or less, and more preferably 0.40% or less.
• the C content is still more preferably 0.35% or less.
• Silicon is added for deoxidization, but the deoxidization effect is insufficient if the content is less than 0.05%.
• the Si content is limited to 0.05% or more.
• the Si content is preferably 0.10% or more, more preferably 0.20% or more, and still more preferably 0.30% or more. If, on the other hand, the Si content is more than 2.00%, the effect is saturated. Thus, the Si content is limited to 2.00% or less.
• the Si content is preferably 1.00% or less, and more preferably 0.80% or less. More than 0.50% silicon deteriorates toughness and weldability. Thus, the Si content is more preferably 0.50% or less.
• manganese is an element that enhances the hardenability of steel and contributes to increasing the strength of steel.
• the Mn content is limited to 0.30% or more.
• the Mn content is preferably 0.40% or more, more preferably 0.45% or more, and still more preferably 0.50% or more.
• manganese is an element that segregates in steel to harden the steel locally. When added in a large amount, manganese forms localized hard regions to disadvantageously lower the hydrogen embrittlement resistance.
• the Mn content in the present invention is limited to 1.50% or less.
• the Mn content is preferably 1.20% or less, more preferably 1.00% or less, and still more preferably 0.80% or less.
• Phosphorus is an element that segregates at grain boundaries in the steel microstructure to cause grain boundary embrittlement, and also segregates to harden the steel locally.
• phosphorus is an incidental impurity. While it is preferable to remove as much phosphorus as possible, up to 0.015% phosphorus is acceptable. Thus, the P content is limited to 0.015% or less.
• the P content is preferably 0.008% or less.
• the P content is more preferably 0.005% or less, and still more preferably 0.003% or less. While a lower content is more preferable, from the point of view of refining costs, the P content is preferably 0.0001% or more, and more preferably 0.001% or more.
• Sulfur is an incidental impurity. In steel, most of sulfur is present as sulfide inclusions that reduce ductility, toughness, and SSC resistance. While it is preferable to remove as much sulfur as possible, up to 0.005% sulfur is acceptable. Thus, the S content is limited to 0.005% or less.
• the S content is preferably 0.003% or less, and more preferably 0.002% or less. While a lower content is more preferable, from the point of view of refining costs, the S content is preferably 0.0002% or more. The S content is more preferably 0.001% or more.
• Aluminum is added as a deoxidizing agent but produces no effects when the amount is less than 0.005%.
• the Al content is limited to 0.005% or more.
• the Al content is preferably 0.010% or more, and more preferably 0.020% or more.
• more than 0.150% aluminum lowers the cleanliness of steel and deteriorates toughness.
• the Al content is limited to 0.150% or less.
• the Al content is preferably 0.130% or less, more preferably 0.100% or less, and most preferably 0.080% or less.
• nitrogen is present as an incidental impurity. Nitrogen has an effect of reducing the size of crystal grains to enhance toughness by combining with aluminum to form AlN, and, when titanium is present, by forming TiN.
• the N content is preferably 0.0005% or more, and more preferably 0.001% or more. If, however, the N content is more than 0.006%, coarse nitrides are formed to cause a significant decrease in toughness. Thus, the N content is limited to 0.006% or less.
• the N content is preferably 0.005% or less, more preferably 0.004% or less, and still more preferably 0.003% or less.
• Chromium is an element that increases the strength of steel by enhancing hardenability and also enhances corrosion resistance. Furthermore, chromium is an element that combines with carbon during tempering to form precipitates, such as M 3 C, M 7 C 3 , and M 23 C 6 (M is the metal element), and enhances temper softening resistance. This element is necessary especially for increasing the strength of steel pipes. In particular, M 3 C precipitates are highly effective in enhancing temper softening resistance. In order to obtain these effects, the Cr content is limited to more than 0.2%. The Cr content is preferably 0.3% or more, and more preferably 0.5% or more. On the other hand, more than 1.7% chromium forms large amounts of M 7 C 3 and M 23 C 6 , which serve as hydrogen trapping sites to reduce hydrogen attack resistance.
• the Cr content is limited to 1.7% or less.
• the Cr content is preferably 1.5% or less, more preferably 1.0% or less, and still more preferably 0.8% or less.
• Molybdenum is an element that forms precipitates and contributes to increasing the strength of steel by precipitation strengthening, and effectively contributes to ensuring the desired high strength while reducing the dislocation density during tempering. Furthermore, molybdenum is dissolved in steel and segregates at prior austenite grain boundaries to contribute to enhancing hydrogen embrittlement resistance. Furthermore, molybdenum has effects of densifying corrosion products and suppressing the generation and growth of pits that serve as the origin of cracks. In order to obtain these effects, the Mo content is limited to more than 1.0%. The Mo content is preferably more than 1.1%, more preferably more than 1.2%, still more preferably 1.3% or more, and most preferably 1.4% or more.
• the Mo content is limited to 3.0% or less.
• the Mo content is preferably 2.8% or less, more preferably 2.5% or less, still more preferably 1.8% or less, and most preferably 1.5% or less.
• Niobium contributes to increasing the strength of steel through precipitation strengthening by forming precipitates or carbonitrides, and also contributes to reducing the size of austenite grains.
• the Nb content is limited to 0.001% or more.
• the Nb content is preferably 0.005% or more, more preferably 0.006% or more, and still more preferably 0.007% or more.
• coarse Nb precipitates tend to serve as the origin of hydrogen-induced cracking and the presence of a large amount of Nb precipitates resulting from the addition of more than 0.020% niobium leads to a significant decrease in hydrogen embrittlement resistance of high-strength steel materials.
• the Nb content in the present invention is limited to 0.020% or less.
• the Nb content is preferably 0.015% or less, and more preferably less than 0.010%.
• the B content is limited to 0.0003% or more.
• the B content is preferably 0.0007% or more, and more preferably 0.0010% or more. If, on the other hand, the B content is more than 0.0030%, boron is precipitated as, for example, carbonitrides to cause a decrease in hardenability and hence a decrease in toughness. Thus, the B content is limited to 0.0030% or less.
• the B content is preferably 0.0025% or less.
• the B content is more preferably 0.0020% or less, and still more preferably 0.0015% or less.
• Oxygen (O) is an incidental impurity and is present as oxide inclusions in steel. These inclusions serve as the origin of in a hydrogen gas environment and lower hydrogen embrittlement resistance. It is therefore preferable in the present invention that oxygen (O) be removed as much as possible. However, excessive deoxidization increases the refining costs. Up to 0.0030% oxygen (O) is acceptable. Thus, the O (oxygen) content is limited to 0.0030% or less. The O content is preferably 0.0025% or less. The O content is more preferably 0.0020% or less. The O content is still more preferably 0.0015% or less. Although the lower limit is not particularly limited, the O content is preferably 0.0010% or more.
• titanium combines with nitrogen and is precipitated as fine TiN, which produces a pinning effect to contribute to reducing the size of austenite grains.
• 0.003% or more titanium needs to be added.
• the effect obtained by titanium is small if the content is less than 0.003%.
• the Ti content is limited to 0.003% or more.
• the Ti content is preferably 0.005% or more, and more preferably 0.010% or more. If, on the other hand, the Ti content is more than 0.025%, TiN is coarsened and fails to produce the pinning effect described above, and further the toughness is deteriorated. Furthermore, the hydrogen embrittlement resistance is lowered due to the coarse TiN. For these reasons, the Ti content is limited to 0.025% or less.
• the Ti content is preferably 0.020% or less, and more preferably 0.015% or less.
• Mo/C is less than 2.0, the amount of Mo precipitates that are formed is reduced due to the lack of molybdenum, and consequently Mo precipitates enough to enhance hydrogen embrittlement resistance are not formed.
• Mo/C is limited to more than 2.0.
• Mo/C is preferably 2.5 or more, more preferably 3.0 or more, and still more preferably 3.5 or more. If, on the other hand, Mo/C is greater than 12.0, Mo precipitates highly tend to become coarse to deteriorate toughness and hydrogen embrittlement resistance. Because part of the coarse Mo precipitates results from the aggregation and coalescence of fine Mo precipitates, the number density of fine Mo precipitates is lowered.
• Mo/C is limited to 12.0 or less.
• Mo/C is preferably 10.0 or less, more preferably 8.0 or less, still more preferably 6.0 or less, and most preferably 5.0 or less.
• the composition may optionally include one, or two or more selected from V: 0.30% or less, Cu: 1.00% or less, Ni: 2.0% or less, and W: 3.0% or less; H: 0.0010% or less; Ca: 0.0005 to 0.005%; or any combination thereof.
• V 0.30% or less
• Cu 1.00% or less
• Ni 2.0% or less
• W 3.0% or less
• Vanadium, copper, nickel, and tungsten are elements that contribute to increasing the strength of steel. One, or two or more may be selectively added as necessary.
• V 0.30% or less
• Vanadium is an element that contributes to strengthening of steel by forming precipitates and carbonitrides.
• the V content may be 0% or more.
• the V content is preferably 0.02% or more, and more preferably 0.03% or more.
• the effect is saturated even when more than 0.30% vanadium is added, and the addition will not produce the corresponding effect and is economically disadvantageous.
• the V content is limited to 0.30% or less.
• the V content is preferably 0.20% or less, and more preferably 0.15% or less.
• Copper is an element that is effective for improving toughness and increasing strength. However, excessive addition deteriorates weldability. Thus, when copper is added, the Cu content is limited to 1.00% or less.
• the Cu content is preferably 0.75% or less, more preferably 0.50% or less, and still more preferably 0.25% or less.
• the Cu content may be 0% or more. In order to obtain the above effects, the Cu content is preferably 0.01% or more.
• Nickel is an element that contributes to increasing the strength of steel and enhances toughness and corrosion resistance.
• the Ni content is desirably 0.03% or more.
• the Ni content is more preferably 0.1% or more.
• the effects are saturated even when more than 2.0% nickel is added, and the addition will not produce the corresponding effect and is economically disadvantageous.
• the Ni content is limited to 2.0% or less.
• the Ni content is preferably 1.5% or less, more preferably 1.0% or less, and still more preferably 0.5% or less.
• the high-strength seamless steel pipe of the present invention has the composition described hereinabove and further has a microstructure in which the main phase is tempered martensite.
• the microstructure contains precipitates, and the precipitates include precipitates having a diameter of 50 nm or less.
• the proportion of molybdenum present in fine precipitates having a diameter of 50 nm or less is limited to 50% by mass or more of the molybdenum present in precipitates.
• the proportion of molybdenum present in fine precipitates having a diameter of 50 nm or less is preferably 60% by mass or more of the molybdenum present in precipitates.
• the proportion is more preferably 65% by mass or more, and still more preferably 70% by mass or more.
• the proportion of molybdenum present in fine precipitates having a diameter of 50 nm or less is preferably 95% by mass or less, more preferably 90% by mass or less, of the molybdenum present in precipitates.
• a billet (a steel pipe material) is formed by a continuous casting method
• the steel being poured from the ladle into the tundish is sealed with an inert gas in order to reduce the amounts of nitride inclusions and oxide inclusions, and electromagnetic stirring is performed in the mold to cause the inclusions to float for separation.
• the refining process is not limited to the above. Appropriate management is important also in other refining processes.
• a steel pipe material having the composition described above is heated and hot-rolled to form a seamless steel pipe with a predetermined shape.
• the seamless steel pipe for high-pressure hydrogen container is preferably applied to a hydrogen container with a hydrogen pressure of 1 MPa or more, more preferably a hydrogen pressure of 20 MPa or more.
• the upper limit of the hydrogen pressure is not particularly limited, but the hydrogen pressure of interest is preferably 120 MPa or less.
• the steel pipe material (hereinafter, also written simply as the steel material) used in the present invention is preferably a billet (a round billet) produced by smelting a molten steel having the above-described composition with a common smelting technique, such as a converter, and casting the steel by a common casting method, such as a continuous casting method.
• the billet may be further hot-rolled into a round bar having a predetermined shape, or a round bar may be produced by ingot making-blooming.
• the steel pipe of the present invention may be manufactured by sequentially performing the following steps (1) to (3).
• the temperature in the following description indicates the temperature on the surface of the billet or the steel pipe.
• the casting speed is preferably 1.8 m/min or less.
• the lower the casting speed the greater the reduction in hydrogen concentration and in the amount of inclusions in the steel. This effect is more marked at the casting speed 1.0 m/min or less.
• the casting speed is more preferably 1.0 m/min or less, still more preferably 0.5 m/min or less, and most preferably 0.1 m/min or less.
• the lower limit is not particularly limited, but the casting speed is preferably 0.01 m/min or more to avoid difficult device control.
• the billet having the above-described chemical composition is heated for hot rolling.
• the billet is not particularly limited and may be, for example, a billet obtained by a usual continuous casting method.
• Heating temperature 1050 to 1350°C
• the heating temperature is limited to 1050°C or above.
• the heating temperature is preferably 1100°C or above, and more preferably 1150°C or above.
• heating at above 1350°C increases the grain size and also results in coarsening of precipitates, such as TiN, that occur during solidification, as well as coarsening of cementite, thereby lowering the toughness of the steel pipe.
• heating to a temperature above 1350°C forms a thick scale layer on the surface of the steel pipe material to cause problems, such as surface defects at the time of rolling, and also increases the energy loss and is disadvantageous from the point of view of energy saving.
• the heating temperature is limited to 1350°C or below.
• the heating temperature is preferably 1300°C or below.
• the heating temperature is more preferably 1250°C or below.
• the billet heated in the heating step is rolled to give a steel pipe with a predetermined shape.
• the rolling may be hot rolling including piercing by a usual Mannesmann-plug mill process or Mannesmann-mandrel mill process.
• the predetermined shape may be, for example, a hollow cylindrical steel pipe shape, with examples including tapered steel pipes with a smaller end diameter than the center diameter, and gas cylinder shapes represented by pressure vessels.
• a hollow cylindrical steel pipe is preferably 200 to 600 mm in outer diameter and 500 to 12000 mm in steel pipe length in the pipe axis direction.
• a tapered steel pipe with a smaller end diameter than the center diameter is preferably 200 to 600 mm in outer diameter of the central portion, 50 to 550 mm in end diameter, and 500 to 12000 mm in steel pipe length in the pipe axis direction.
• a gas cylinder-shaped steel pipe is preferably 200 to 600 mm in outer diameter and 500 to 12000 mm in cylinder length in the pipe axis direction.
• the hot rolling also includes a process that performs a step of rolling a billet into a steel pipe shape (a hot working step) and an expansion step simultaneously. Where necessary, a sizing step for adjusting the wall thickness may be carried out after the reheating step described later.
• the seamless steel pipe obtained is subjected to a cooling treatment in which the seamless steel pipe is cooled at a cooling rate equal to or faster than air cooling until the surface temperature reaches 200°C or below.
• Cooling treatment after the completion of hot rolling average cooling rate: equal to or faster than air cooling, cooling stop temperature: 200°C or below
• the chemical composition of the present invention can give a microstructure containing a martensite phase as the main phase when the steel pipe from the hot rolling is cooled at an average cooling rate equal to or faster than air cooling. If the air cooling (the cooling) is stopped when the surface temperature is still above 200°C, the transformation may not be fully completed. Thus, the cooling treatment after the hot rolling is to be performed at an average cooling rate equal to or faster than air cooling until the surface temperature reaches 200°C or below.
• the "cooling rate equal to or faster than air cooling” indicates 0.1°C/s or more. If the average cooling rate is less than 0.1°C/s, the metallic microstructure after the cooling becomes nonuniform, and the subsequent heat treatment gives a nonuniform metallic microstructure.
• the average cooling rate is preferably 1.0°C/s or more, and more preferably 10.0°C/s or more. Although the upper limit is not particularly limited, the average cooling rate is preferably 1000.0°C/s or less.
• the average cooling rate is the average of the cooling rates from Ac 3 transformation temperature to 200°C.
• Reheating temperature for quenching equal to or higher than Ac 3 transformation temperature and equal to or lower than 1000°C
• the reheating temperature for performing quenching is below Ac 3 transformation temperature, the steel is not heated to the austenite single phase region and consequently the microstructure that is obtained will not contain a martensite phase as the main phase.
• the reheating temperature is limited to equal to or higher than Ac 3 transformation temperature.
• the reheating temperature is preferably Ac 3 temperature + 30°C or above, and more preferably Ac 3 temperature + 50°C or above.
• the reheating temperature is not set to Ac 3 temperature + 30°C or above or to Ac 3 temperature + 50°C or above when "Ac 3 temperature + 30°C" or "Ac 3 temperature + 50°C" for the system exceeds 1000°C.
• reheating at above 1000°C is disadvantageous in that, for example, grains are coarsened to cause a decrease in toughness, and oxide scales on the surface are increased in thickness and easily come off to cause defects on the steel sheet surface. Furthermore, such reheating applies an excessively high load to the heat treatment furnace and is problematic from the point of view of energy saving.
• the reheating temperature for quenching is limited to 1000°C or below, and is preferably 980°C or below, and more preferably 950°C or below.
• the reheated steel is quenched.
• the cooling in the quenching treatment is effected by rapidly cooling the steel until the surface temperature is lowered to 200°C or below.
• quenching or rapid cooling means that the average cooling rate from Ac 3 transformation temperature to 200°C is 2.0°C/s or more.
• the average cooling rate is preferably 5.0°C/s or more, and more preferably 10.0°C/s or more.
• the upper limit is not particularly limited, but the average cooling rate is preferably 1000.0°C/s or less.
• water cooling preferably lowers the temperature at the center of the wall thickness from Ac 3 transformation temperature to 400°C or below at an average cooling rate of 2.0°C/s or more.
• the upper limit is not particularly limited, but this average cooling rate is preferably 1000.0°C/s or less.
• the cooling is preferably effected until the surface temperature is lowered to 100°C or below.
• the surface temperature after the cooling is preferably low and is preferably room temperature.
• the quenching treatment may be repeated two or more times.
• the Ac 3 transformation temperature used here is a value calculated from the following expression.
• Ac 3 transformation temperature (°C) 937 - 476.5C + 56Si - 19.7Mn - 16.3Cu - 4.9Cr - 26.6Ni + 38.1Mo + 124.8V + 136.3Ti + 198Al + 3315B (Here, C, Si, Mn, Cu, Cr, Ni, Mo, V, Ti, Al, and B indicate the contents (mass%) of the respective elements.) When any of the elements described in the above expression is not contained in the steel, the content of that element is taken as 0% in the calculation of Ac 3 transformation temperature.
• the steel After cooled at a cooling rate equal to or faster than air cooling, the steel is tempered.
• the tempering treatment heats the steel to a temperature in the range of 600 to 740°C.
• Tempering temperature 600 to 740°C
• Tempering is performed for the purposes of reducing the dislocation density and precipitating Mo precipitates, thereby enhancing toughness and hydrogen embrittlement resistance.
• the reduction in dislocation and the precipitation of Mo precipitates are insufficient at a tempering temperature of below 600°C, and the treatment fails to ensure excellent hydrogen embrittlement resistance.
• the tempering temperature is limited to 600°C or above.
• the tempering temperature is preferably 620°C or above, more preferably 640°C or above, and still more preferably 660°C or above.
• tempering at a temperature above 740°C significantly softens the microstructure and the desired high strength cannot be ensured.
• the tempering temperature is limited to 740°C or below.
• the tempering temperature is preferably 710°C or below.
• the tempering temperature is more preferably 700°C or below, and still more preferably 680°C or below.
• Average heating rate until the tempering temperature is reached 0.5°C/min or more
• the average heating rate until the tempering temperature is reached is limited to 0.5°C/min or more, and is preferably 1.0°C/min or more, more preferably 2.0°C/min or more, and most preferably 5.0°C/min or more. There is no particular upper limit, but excessively rapid heating produces nonuniform temperature distribution and results in inhomogeneous material microstructure.
• the average heating rate is preferably 50.0°C/min or less.
• Holding time at the tempering temperature 10 minutes or more and less than 60 minutes
• the holding time at the tempering temperature is limited to 10 minutes or more.
• the holding time at the tempering temperature is preferably 15 minutes or more, and more preferably 20 minutes or more. If the holding time at the tempering temperature is too long, the size of precipitates is excessively increased. Thus, the holding time is limited to less than 60 minutes. Because holding time is a cost increasing factor in terms of energy, the tempering time is preferably less than 50 minutes, more preferably less than 40 minutes, and still more preferably less than 30 minutes.
• the steel after being hot-rolled is subjected to the cooling treatment in which the steel is cooled at a cooling rate equal to or faster than air cooling, and is reheated and quenched at least one time by water cooling or the like, and is tempered in the manner described above.
• the quenching be performed five times or less.
• a warm or cold straightening treatment may be performed as required to correct shape defects of the steel pipe.
• Table 2 describes the tempering conditions, the tempered martensite area fraction, the prior austenite grain size index, the proportion by mass% of molybdenum contained in precipitates relative to the molybdenum contained in the steel, the proportion of molybdenum contained in precipitates having a diameter of 50 nm or less relative to the molybdenum contained in precipitates, TS, and the relative reduction of area (RRA) for each of Nos. 1 to 24.
• Billets having the chemical compositions described in Nos. 1 to 24 in Table 1 were produced at a casting speed of 0.6 m/min.
• the billets were heated to 1250°C, hot-worked, and expanded to give seamless steel pipes.
• the conditions in the production of the seamless steel pipes were such that the expansion was completed at 820°C or above, and the hot-worked pipes were cooled at a cooling rate equal to or faster than air cooling to such a temperature that the surface temperature was 200°C or below.
• Those steel pipes having an Ac 3 transformation temperature of 950°C or below were heated and held at 950°C, and those steel pipes having an Ac 3 transformation temperature of above 950°C were heated and held at 1000°C.
• the steel pipes were then water-cooled at 5.0°C/s to 200°C or below, and were subsequently tempered.
• the tempering step was performed at the heating rate, the holding temperature, and the holding time described in Table 2.
• the tempering temperature was controlled so that the tensile strength would be in the range of 850 to 950 MPa.
• the steel pipes obtained were analyzed to evaluate the metallic microstructure and mechanical properties.
• steel pipes Nos. 25 to 39 in Table 3 were manufactured as follows. Any of the billets Nos. 5, 8, and 12 having the chemical compositions described in Table 1 were produced at various casting speeds. The billets were heated to 1250°C and expanded to give seamless steel pipes. The conditions in the steel pipe production were such that the expansion was completed at 820°C or above, and the hot-worked pipes were cooled at a cooling rate equal to or faster than air cooling to such a temperature that the surface temperature was 200°C or below. Those steel pipes having an Ac 3 transformation temperature of 950°C or below were heated and held at 950°C, and those steel pipes having an Ac 3 transformation temperature of above 950°C were heated and held at 1000°C. The steel pipes were then water-cooled at 5.0°C/s to 200°C or below, and were subsequently tempered under conditions described in Table 3. The steel pipes obtained were analyzed to evaluate the metallic microstructure and mechanical properties.
• the evaluation methods are as follows.
• the metallic microstructure at 1/4 wall thickness on the inner side of the steel pipe was evaluated as follows. Samples were obtained from a cross section parallel to the longitudinal direction and the wall thickness direction of the steel pipe, so that positions at 1/4 wall thickness on the inner side and at the center of the wall thickness would be the observation faces. The cross sections of the samples were etched with a 3 vol% Nital solution and were photographed with a scanning electron microscope at an appropriate magnification between 1000 and 5000 times. Tempered martensite, ferrite, bainite, and pearlite were observed. Tempered martensite was identified visually by comparing the image with the microstructure image shown in Reference 2.
• the SEM image was divided into regions based on the above identification, and the resultant image was binarized by image analysis into martensite and other regions.
• the fraction of tempered martensite was determined as the area fraction of tempered martensite.
• the steel material taken from the steel pipe was analyzed as follows to measure Mo precipitates. Mo precipitates were identified by an extraction method in which the steel material was electrolyzed and the resultant precipitates were filtered.
• a 10 mm square sample taken from a cross section perpendicular to the rolling direction of the steel pipe (a cross section perpendicular to the pipe axis direction: C cross section) was electrolyzed at a constant current in a 10% AA electrolytic solution to dissolve the steel.
• the residue was added to a 0.05 wt% aqueous sodium hexametaphosphate solution and was ultrasonicated. Precipitates were thus isolated.
• the dispersion was filtered through a 50 nm mesh filter, and thereby 50 nm and smaller precipitates were obtained.
• the 50 nm and smaller precipitates that had passed through the filter, and more than 50 nm precipitates remaining on the filter were each thermally treated with white fuming sulfuric acid, perchloric acid, and nitric acid, and were then dissolved into hydrochloric acid.
• the precipitate solutions and the electrolytic solution containing the dissolved portion were each analyzed by ICP to determine the Mo concentration (mass%) in the precipitates of each size group and the concentration (mass%) of dissolved molybdenum.
• the amounts of molybdenum contained in all the precipitates, and the amount of dissolved molybdenum obtained as described above were combined to give the total amount of molybdenum contained in the steel.
• test piece for tensile test was taken from a cross section perpendicular to the steel pipe axis (C direction) in such a manner that a position at 1/4 wall thickness on the inner side of the steel pipe would be the center of the test piece and the longitudinal direction of the test piece would be the C direction.
• the test piece used here was a bar-shaped test piece specified in JIS Z 2201 "Tensile test pieces for metallic materials". The test was performed using the method specified in JIS Z2241, and the maximum load was taken as TS of the steel pipe. While it is preferable that the sampling is centered at 1/4 wall thickness, the center of sampling may be other than at 1/4 wall thickness when the steel pipe has a small wall thickness (for example, a wall thickness of 45 mm or less).
• Hydrogen embrittlement resistance was evaluated based on the relative reduction of area (RRA) of a test specimen after a slow strain rate tensile test in hydrogen gas in accordance with ASTM G 142.
• RRA relative reduction of area
• a steel material undergoes plastic deformation and the fracture surface has a small area, and thus the reduction of area ⁇ air is large.
• a steel material exhibits less elongation and is fractured before the area is reduced, and consequently the fracture surface has a large area.
• the reduction of area ⁇ H of the fracture surface after the test in hydrogen is small, unlike the testing in air.
• Hydrogen embrittlement resistance was evaluated based on how much the reduction of area had decreased.
• the relative reduction of area obtained by a slow strain rate tensile test (stress rate: 0.002 mm/s) at room temperature under 105 MPa hydrogen gas is described in Table 2. The larger the RRA, the higher the hydrogen embrittlement resistance. In this evaluation, 60% or higher RRA was accepted.
• ⁇ air is the ratio "sectional area of the test specimen after testing in air/sectional area before testing”
• ⁇ H is the ratio "sectional area of the test specimen after testing in hydrogen/sectional area before testing”.

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